CN1557006A - Method for producing R-T-B based rare earth permanent magnet - Google Patents

Abstract

In order to obtain a sintered body having a composition of 25-35 wt% R (R refers to one or more rare earth elements, but the term "rare earth elements" refers to a concept including Y), 0.5-4.5 wt% B, 0.02-0.6 wt% of one or two of Al and Cu, 0.03-0.25 wt% Zr, 4 wt% or less Co (excluding 0), and the remainder being essentially Fe, with a coefficient of variation (CV value) indicating Zr dispersibility of 130 or less, Zr is incorporated into a low-R alloy when producing an R-T-B rare earth permanent magnet using a mixing method. This sintered body minimizes degradation of magnetic properties, suppresses grain growth, and improves the sintering temperature range.

Description

Manufacturing method of R-T-B series rare earth permanent magnet

technical field

The present invention relates to R (R is one or more than two kinds of rare earth elements, but the concept of rare earth elements containing Y), T (at least one or more transition metal elements that are necessary for Fe or Fe and Co) And the manufacturing method of the R-T-B system rare-earth permanent magnet mainly composed of B (boron).

Background technique

Among the rare-earth permanent magnets, the R-T-B series rare-earth permanent magnets have excellent magnetic properties, and the main component Nd is abundant and relatively cheap, so the demand is increasing year by year.

Research and development to improve the magnetic properties of R-T-B series rare earth permanent magnets are being vigorously carried out. For example, Japanese Unexamined Patent Publication No. 1-219143 reports that adding 0.02 to 0.5 atomic % of Cu to an R-T-B-based rare earth permanent magnet can improve magnetic properties and improve heat treatment conditions. However, the method described in JP-A-1-219143 is not sufficient for obtaining high magnetic properties required for high-performance magnets, specifically for obtaining relatively high coercive force (HcJ) and residual magnetic flux density (Br). full.

Here, the magnetic properties of the R-T-B based rare earth permanent magnet obtained by sintering may depend on the sintering temperature. On the other hand, on an industrial scale, it is difficult to make the heating temperature uniform over the entire area in the sintering furnace. Therefore, for R-T-B series rare earth permanent magnets, it is required to obtain desired magnetic properties even if the sintering temperature fluctuates. Here, the temperature range in which desired magnetic properties can be obtained is called the sintering temperature range.

In order to make the R-T-B series rare earth permanent magnet into a higher-performance permanent magnet, it is necessary to reduce the oxygen content in the alloy. However, when the oxygen content in the alloy is reduced, abnormal grain growth tends to occur in the sintering process, and the squareness ratio (also referred to as squareness ratio) decreases. This is because the oxide formed by the oxygen in the alloy suppresses the grain growth.

Here, a method of adding a new element to a Cu-containing R-T-B-based rare-earth permanent magnet is considered as a means for improving magnetic properties. It is reported in JP-A-2000-234151 that Zr and/or Cr are added in order to obtain high coercive force and residual magnetic flux density.

Similarly, it is reported in JP-A-2002-75717 that by making the fine ZrB compound, NbB compound or HfB compound (hereinafter referred to as M-B Compounds) are evenly dispersed and precipitated, inhibiting the grain growth in the sintering process, and improving the magnetic properties and sintering temperature range.

According to Japanese Patent Application Laid-Open No. 2002-75717, the sintering temperature range can be widened by dispersing and depositing the M-B compound. However, in Example 3-1 disclosed in Japanese Unexamined Patent Publication No. 2002-75717, the sintering temperature range is about 20° C. which is relatively narrow. Therefore, in order to obtain high magnetic properties in mass-produced sintering furnaces, etc., it is desired to widen the sintering temperature range again. Also, in order to obtain a sufficiently wide sintering temperature range, it is effective to increase the amount of Zr added. However, as the amount of Zr added increases, the remanence magnetic flux density decreases, and the original high characteristics cannot be obtained.

Contents of the invention

Therefore, an object of the present invention is to provide a method for producing an R-T-B based rare earth permanent magnet capable of minimizing deterioration of magnetic properties, suppressing growth of crystal grains, and further improving the sintering temperature range.

In recent years, in the production of high-performance RTB-based rare earth permanent magnets, mixing and sintering various metal powders and alloy powders of different compositions has become mainstream. This mixing method is typically R ₂ T ₁₄ B-based intermetallic compound (R is one or more than two kinds of rare earth elements (but the concept of rare earth elements contains Y), T is Fe or Fe and Co are At least one transition metal element as the main body) The alloy for forming the main phase is mixed with the alloy for forming the grain boundary phase existing between the main phases (hereinafter referred to as "the alloy for forming the grain boundary phase") . Here, since the content of the rare earth element R in the alloy forming the main phase is relatively small, it may be called a low-R alloy. On the other hand, alloys for forming grain boundaries are sometimes called high-R alloys because they contain a relatively large amount of rare earth element R.

The present inventors have confirmed that when the R-T-B system rare earth permanent magnet is obtained by the mixing method, Zr dispersibility in the obtained R-T-B system rare earth permanent magnet is high when the low R alloy is made to contain Zr. Due to the high dispersion of Zr, it is possible to prevent abnormal grain growth and further expand the sintering temperature range with a lower Zr content.

The present invention provides a method for manufacturing an RTB system rare earth permanent magnet based on the above knowledge, wherein the RTB system rare earth permanent magnet is made of a sintered body with the following composition: R: 25 to 35% by weight (R One or two or more rare earth elements, but the rare earth element is the concept of containing Y), B: 0.5 to 4.5% by weight, one or two of Al and Cu: 0.02 to 0.6% by weight, Zr: 0.03 ～0.25% by weight, Co: 4% by weight or less (excluding 0), and the remainder is substantially composed of Fe. The method includes: making a low R alloy containing Zr mainly composed of R ₂ T ₁₄ B ₁ compounds, And a formed body containing a high R alloy mainly composed of R and T, and the formed body is sintered.

In this production method, it is preferable to add Zr to the low-R alloy, and to further contain one or both of Cu and Al. This is because it is effective to improve the dispersibility of Zr in the low R alloy by including one or both of Cu and Al.

As just explained, according to the R-T-B series rare earth permanent magnet of the present invention, the sintering temperature range can be improved. The effect of improving the sintering temperature range depends on the magnet composition in the state of powder (or its molded body) before sintering. Therefore, in the molded body of the present invention, the sintering temperature range at which the square ratio (Hk/HcJ) of the R-T-B series rare earth permanent magnet obtained by sintering is 90% or more is 40° C. or more.

For the R-T-B series rare earth permanent magnet of the present invention, Zr is preferably 0.05-0.2% by weight, more preferably 0.1-0.15% by weight.

Also, for the R-T-B series rare earth permanent magnet of the present invention, as a composition other than Zr, R: 28 to 33% by weight, B: 0.5 to 1.5% by weight, Al: 0.30% by weight or less (excluding 0), Cu: 0.3% by weight or less (not including 0), Co: 0.1 to 2.0% by weight or less, and the remainder substantially composed of Fe is ideal, consisting of R: 29 to 32% by weight, B: 0.8 to 1.2% by weight, Al: 0.25% by weight or less (excluding 0), Cu: 0.15% by weight or less (excluding 0), and a composition in which the remainder is substantially composed of Fe is also preferable.

In addition, the effect of improving the dispersibility of Zr and expanding the sintering temperature range by adding Zr to the low R alloy is remarkable when the amount of oxygen contained in the sintered body is low oxygen content of 2000 ppm or less.

Description of drawings

FIG. 1 is a graph showing chemical compositions of low-R alloys and high-R alloys used in the first embodiment.

2 is a graph showing the final composition, oxygen content, and magnetic properties of permanent magnets (No. 1 to 20) obtained in the first example.

3 is a graph showing the final composition, oxygen content, and magnetic properties of the permanent magnets (No. 21 to 35) obtained in the first example.

Fig. 4 is a graph showing the relationship between the residual magnetic flux density (Br), coercive force (HcJ) and square ratio (Hk/HcJ) of the permanent magnet (sintering temperature at 1070°C) obtained in the first example and the amount of Zr added .

Fig. 5 is a graph showing the relationship between the residual magnetic flux density (Br), coercive force (HcJ) and square ratio (Hk/HcJ) of the permanent magnet (sintering temperature at 1050°C) obtained in the first example and the amount of Zr added .

6 is a photograph showing the results of EPMA (Electron Probe Micro Analyzer: Electron Probe Micro Analyzer) elemental mapping of the permanent magnet (high-R alloy-added permanent magnet) obtained in the first example.

Fig. 7 is a photograph showing the results of EPMA elemental mapping of the permanent magnet (low-R alloy-added permanent magnet) obtained in the first example.

8 is a graph showing the relationship between the method of adding Zr, the amount of Zr added, and the CV value (coefficient of variation) of Zr in the permanent magnet obtained in the first example.

9 is a graph showing the final composition, oxygen content, and magnetic properties of permanent magnets (No. 36 to 75) obtained in Example 2. FIG.

Fig. 10 is a graph showing the relationship between remanence (Br), coercive force (HcJ) and squareness ratio (Hk/HcJ) and Zr addition amount in the second embodiment.

Fig. 11 is a photograph of the cross-section of permanent magnets No. 37, No. 39, No. 43, and No. 48 obtained in Example 2 observed by SEM (scanning electron microscope).

Fig. 12 is a graph showing 4πI-H curves of the permanent magnets No. 37, No. 39, No. 43 and No. 48 obtained in the second example.

13 is a photograph showing scanned images (30 μm×30 μm) of elements such as B, Al, Cu, Zr, Co, Nd, Fe, and Pr of the No. 70 permanent magnet obtained in the second example.

Fig. 14 is a diagram showing an example of a profile of EPMA line analysis of the No. 70 permanent magnet obtained in the second example.

Fig. 15 is another example of a profile of the EPMA line analysis of the No. 70 permanent magnet obtained in the second example.

Fig. 16 is a graph showing the relationship between the addition amount of Zr and the sintering temperature and the squareness ratio (Hk/HcJ) in the second embodiment.

Fig. 17 is a graph showing the final composition, oxygen content and magnetic properties of the permanent magnets (No. 76 to 79) obtained in the third example.

18 is a graph showing the final composition, oxygen content, and magnetic properties of permanent magnets (No. 80 to 81) obtained in the fourth example.

Detailed ways

Hereinafter, embodiments of the present invention will be described.

<organization>

First, the structure of the R-T-B based rare earth permanent magnet which is the characteristic of the present invention will be described.

The R-T-B series rare earth permanent magnet of the present invention is characterized in that Zr is uniformly dispersed in the structure of the sintered body. This feature is more specifically designated by the coefficient of variation (referred to as CV (Coefficient of Variation) in the specification of this application; also referred to as the coefficient of variation). In the present invention, the CV value of Zr is 130 or less, preferably 100 or less, more preferably 90 or less. The smaller the CV value, the higher the degree of dispersion of Zr. Also, as is well known, the CV value is a quotient (percentage) obtained by dividing the standard deviation by the arithmetic mean. In addition, the CV value of this invention is the value calculated|required from the measurement conditions of the Example mentioned later.

Thus, the high dispersibility of Zr is attributed to the addition method of Zr. As will be described later, the R-T-B based rare earth permanent magnet of the present invention can be produced by a hybrid method. The mixing method is to mix the low-R alloy for forming the main phase with the high-R alloy for forming the grain boundary phase. When the low-R alloy contains Zr, its dispersibility is remarkable compared with the case where the high-R alloy contains Zr. improve.

The R-T-B series rare earth permanent magnet of the present invention has a high dispersion degree of Zr, so even if a small amount of Zr is added, the effect of inhibiting grain growth can still be exerted.

Secondly, it can be confirmed that the R-T-B series rare earth permanent magnet of the present invention: ① Cu and Co can be enriched in the Zr-enriched area at the same time, ② Cu and Co can be enriched in the Zr-enriched area at the same time, ③ Can be enriched in the Zr-enriched area at the same time Cu, Co and Nd. In particular, the ratio of Zr and Cu co-enrichment is high, and Zr and Cu co-exist to exert its effect. In addition, Nd, Co, and Cu are all elements that form grain boundary phases. Therefore, since Zr is enriched in its region, it can be judged that Zr exists in the grain boundary phase.

The reason why Zr, Cu, Co, and Nd exhibit the above-mentioned existence form is not conclusive, but it is considered as follows.

According to the present invention, a liquid phase in which one or more of Cu, Nd, and Co is enriched together with Zr (hereinafter referred to as "Zr-enriched liquid phase") is generated during the sintering process. This Zr-rich liquid phase has a different wettability to the R ₂ T ₁₄ B ₁ crystal grains (compounds) than the usual Zr-free liquid phase. This is a factor that inactivates the grain growth rate during sintering. It is therefore possible to suppress the growth of crystal grains and prevent the generation of huge abnormal crystal grains. At the same time, since the Zr-rich liquid phase may improve the sintering temperature range, it is possible to easily manufacture an RTB-based rare earth permanent magnet with high magnetic characteristics.

When one or more of Cu, Nd, and Co form an enriched grain boundary phase together with Zr, the above effects can be obtained. Therefore, it is possible to distribute them uniformly and finely compared to the case where they exist in a solid state during sintering (oxides, borides, etc.). From this, it is presumed that the necessary addition amount of Zr can be reduced without causing a large amount of out-of-phase such as reducing the main phase ratio, so that the reduction of magnetic properties such as residual magnetic flux density (Br) will not be caused.

<chemical composition>

Next, the ideal chemical composition of the R-T-B based rare earth permanent magnet of the present invention will be described. The chemical composition mentioned here refers to the chemical composition after sintering. As will be described later, the R-T-B based rare earth permanent magnet of the present invention can be produced by the hybrid method, and the various alloys of low R alloy and high R alloy used by the hybrid method will be mentioned in the description of the production method.

The R-T-B series rare earth permanent magnet of the present invention contains 25 to 35% by weight of R.

Here, R is one or more selected from La, Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho, Er, Yb, Lu, and Y. When the amount of R is less than 25% by weight, the formation of the R ₂ T ₁₄ B ₁ phase which is the main phase of the rare earth permanent magnet is insufficient. Therefore, α-Fe with soft magnetic properties is precipitated, _and the _coercive force _is significantly reduced; Density drops. When the amount of R exceeds 35% by weight, R reacts with oxygen, and the amount of oxygen contained increases, and the R-rich phase effective for generating coercive force decreases, resulting in a decrease in coercive force. Therefore, the amount of R is determined to be 25 to 35% by weight. The ideal R amount is 28 to 33% by weight, and the more desirable R amount is 29 to 32% by weight.

Nd is rich in resources and relatively cheap, so it is ideal to choose Nd as the main component of R. In addition, it is effective to contain Dy in order to increase the anisotropic magnetic field and improve the coercive force. Therefore, Nd and Dy are selected as R, and the total amount of Nd and Dy is preferably 25 to 33% by weight. Furthermore, the amount of Dy within this range is preferably 0.1 to 8% by weight. It is preferable to determine the amount of Dy within the above-mentioned range according to the respective degrees of emphasis on the remanence magnetic flux density and the coercive force. That is, when a high residual magnetic flux density is desired, the amount of Dy is 0.1 to 3.5% by weight; when a high coercive force is desired, the amount of Dy is preferably 3.5 to 8% by weight.

Moreover, the R-T-B series rare earth permanent magnet of the present invention contains boron (B) in an amount of 0.5 to 4.5% by weight. When B is less than 0.5% by weight, a high coercive force cannot be obtained; however, when B exceeds 4.5% by weight, the residual magnetic flux density tends to decrease. Therefore, the upper limit is 4.5% by weight. The desirable B content is 0.15 to 1.5% by weight, and the more desirable B content is 0.8 to 1.2% by weight.

The R-T-B based rare earth permanent magnet of the present invention can contain one or both of Al and Cu in the range of 0.02 to 0.6% by weight. By containing one or both of Al and Cu within this range, it becomes possible to increase the coercive force, increase the corrosion resistance, and improve the temperature characteristics of the obtained permanent magnet. When adding Al, the desirable amount of Al is 0.03 to 0.3% by weight, and the more desirable amount of Al is 0.05 to 0.25% by weight. Also, when Cu is added, the amount of Cu is 0.3% by weight or less (excluding 0), preferably 0.15% by weight or less (excluding 0), and more preferably 0.03 to 0.08% by weight.

The R-T-B series rare earth permanent magnet of the present invention contains 0.03 to 0.25% by weight of Zr. In order to improve the magnetic properties of R-T-B series rare earth permanent magnets, Zr exerts the effect of inhibiting the abnormal growth of grains in the sintering process when the oxygen content is reduced, and makes the structure of the sintered body uniform and fine. Therefore, the effect of Zr is remarkable when the oxygen content is low. The preferable content of Zr is 0.05 to 0.2 weight%, and the more preferable content is 0.1 to 0.15 weight%.

The oxygen content of the R-T-B series rare earth permanent magnet of the present invention is below 2000ppm. When the oxygen content is high, the oxide phase which is a non-magnetic component increases, and the magnetic properties are lowered. Here, in the present invention, the oxygen content in the sintered body is determined to be 2000 ppm or less, preferably 1500 ppm or less, more preferably 1000 ppm or less. However, simply reducing the oxygen content reduces the oxide phase which has the effect of suppressing grain growth, and tends to cause grain growth in the process of obtaining a sufficient increase in density during sintering. Here, in the present invention, the R-T-B series rare earth permanent magnet contains Zr in a predetermined amount, which exhibits the effect of suppressing abnormal grain growth during sintering.

The R-T-B series rare earth permanent magnet of the present invention contains Co below 4% by weight (excluding 0), preferably 0.1-2.0% by weight of Co, more preferably 0.3-1.0% by weight. Co forms the same phase as Fe, and is effective in raising the Curie temperature and improving the corrosion resistance of the grain boundary phase.

<Manufacturing method>

Next, preferred embodiments of the method for producing the R-T-B-based rare earth permanent magnet of the present invention will be described.

The present invention uses an alloy mainly containing R ₂ T ₁₄ B phase (low R alloy) and an alloy containing more R than the low R alloy (high R alloy) to manufacture RTB-based rare earth permanent magnets.

First, low R alloys and high R alloys are obtained by performing strip casting of raw metals in vacuum or inert gas, preferably in an Ar gas protective atmosphere. As the raw material metal, rare earth metals or rare earth alloys, pure iron, ferroboron, alloys thereof, and the like can be used. When solidification segregation exists in the obtained master alloy, solution treatment is performed as necessary. The condition is to keep the temperature in the range of 700-1500° C. for more than 1 hour in vacuum or under the protection atmosphere of Ar gas.

The characteristic feature of the present invention is that Zr is added to the low R alloy. This is because the dispersibility of Zr in the sintered body can be improved by adding Zr to the low R alloy as described in the column of <Structure>.

In addition to R, T, and B, the low-R alloy can contain Cu and Al. At this time, the low R alloy constitutes an alloy of the R—Cu—Al—Zr—T(Fe)—B system. In addition, in addition to R, T (Fe) and B, Cu, Co and Al can be contained in the high R alloy. At this time, the high R alloy constitutes an alloy of the R-Cu-Co-Al-T(Fe-Co)-B system.

After the low-R alloy and the high-R alloy are produced, their parent alloys can be pulverized separately or together. The pulverization process includes a coarse pulverization process and a fine pulverization process. First, each master alloy is roughly pulverized to a particle diameter of several hundreds of μm. Coarse pulverization is preferably carried out in an inert gas atmosphere using a pounder, a jaw crusher, a Brown mill, or the like. In order to improve the coarse pulverization property, it is effective to perform coarse pulverization after absorbing hydrogen. In addition, after the hydrogen is stored, the coarse pulverization may be performed after the hydrogen is released.

After the coarse pulverization process is completed, it moves to the fine pulverization process. Fine pulverization mainly uses a jet mill, and the coarse powder with a particle diameter of about several hundred μm is crushed to an average particle diameter of 3 to 5 μm. The jet mill emits high-pressure inert gas (such as nitrogen) from a narrow nozzle to generate a high-speed gas flow, and the high-speed gas flow accelerates the coarsely pulverized powder so that the coarsely pulverized powder collides with each other and the target. Or the method of crushing by the impact of the container wall.

In the fine pulverization step, when the low R alloy and the high R alloy are pulverized separately, the finely pulverized low R alloy powder and the high R alloy powder are mixed in a nitrogen atmosphere. The mixing ratio of the low R alloy powder and the high R alloy powder may be about 80:20 to 97:3 by weight. Similarly, when the low R alloy powder and the high R alloy powder are pulverized together, the mixing ratio may be about 80:20 to 97:3 by weight. When finely pulverizing, by adding additives such as zinc stearate at about 0.01 to 0.3% by weight, fine powder with high orientation can be obtained during molding.

Next, the mixed powder composed of low R alloy powder and high R alloy powder is filled into a mold surrounded by electromagnets, and a magnetic field is applied to make the crystal axes in an oriented state and formed in the magnetic field. Molding in this magnetic field may be carried out in a magnetic field of 12.0 to 17.0 kOe at a pressure of about 0.7 to 1.5 t/cm ² .

After forming in a magnetic field, the formed body is sintered in vacuum or in an inert gas atmosphere. It is necessary to adjust the sintering temperature according to various conditions such as composition, crushing method, particle size and particle size distribution, and sintering at 1000-1100°C for about 1-5 hours is sufficient.

After sintering, aging treatment may be applied to the obtained sintered body. Aging treatment is important in controlling the coercive force. When the aging treatment is carried out in two stages, it is effective to hold the heat at around 600°C and around 800°C for a given period of time. The coercivity increases when heat treatment near 800°C is performed after sintering, so the mixing method is particularly effective. Also, since heat treatment at around 600°C greatly increases the coercive force, when aging treatment is performed in one stage, aging treatment at around 600°C is sufficient.

The rare earth permanent magnet of the present invention with the above composition and manufacturing method can achieve high performance with a remanence (Br) and a coercive force (HcJ) of Br+0.1×HcJ of 15.2 or more, further 15.4 or more.

(Example)

Hereinafter, the present invention will be described in more detail with reference to specific examples. In addition, the R-T-B rare earth permanent magnet of the present invention will be described below for the first to fourth embodiments. The prepared raw material alloy and each manufacturing process have common points, so this point will be described first.

1) Raw material alloy

The 13 alloys shown in Fig. 1 were produced by strip casting.

2) Hydrogen crushing process

After allowing hydrogen to be absorbed at room temperature, dehydrogenation was performed in an Ar atmosphere at 600° C. for 1 hour, followed by hydrogen pulverization.

In order to obtain high magnetic properties, in order to control the oxygen content of the sintered body below 2000ppm in this test, the protective atmosphere of each process from hydrogen treatment (recovery after crushing treatment) to sintering (putting into the sintering furnace) was controlled to less than 100ppm oxygen concentration. Hereinafter referred to as anaerobic process.

3) Crushing process

Generally, two-stage pulverization of coarse pulverization and fine pulverization is carried out. Since the coarse pulverization process cannot be performed under an oxygen-free process, the coarse pulverization process is omitted in this embodiment.

Additives are mixed prior to fine grinding. The types of additives are not particularly limited, as long as they are properly selected to improve the pulverization and the orientation during molding. In this embodiment, 0.05-0.1% of zinc stearate is mixed. The mixing of the additives may be performed, for example, in a Nauta mixer (also referred to as a screw mixer) for about 5 to 30 minutes.

Then, it is finely pulverized by a jet mill until the alloy powder has an average diameter of about 3 to 6 μm. In this experiment, two kinds of pulverized powders having an average particle diameter of 4 μm and 5 μm were produced.

Of course, the additive mixing process and fine pulverization process are all carried out under anaerobic process.

4) Coordination process

In order to conduct experiments efficiently, sometimes several types of fine powders are blended and mixed so as to obtain the desired composition (especially the amount of Zr). The mixing at this time may be performed for about 5 to 30 minutes, for example, with a Nauta mixer or the like.

Although it is ideal to carry out in an oxygen-free process, when the oxygen content of the sintered body is slightly increased, the oxygen content of the fine powder for molding is adjusted by means of this step. For example, a fine powder having the same composition as the average particle diameter is prepared and left in an atmosphere containing oxygen of more than 100 ppm for several minutes to several hours to obtain a fine powder containing several thousand ppm of oxygen. The two kinds of fine powders are mixed in an oxygen-free process to adjust the oxygen content. First Embodiment Various kinds of permanent magnets are produced according to the method described above.

5) Forming process

The resulting fine powder is shaped in a magnetic field. Specifically, the fine powder is filled into a mold surrounded by electromagnets, and the crystal axes are oriented in a magnetic field by applying a magnetic field. Molding in this magnetic field can be done in a magnetic field of 12.0 to 17.0 kOe with a pressure of about 0.7 to 1.5 t/cm ² . In this experiment, molding was performed at a pressure of 1.2t/cm ² in a magnetic field of 15kOe to obtain a molded body. This process is also carried out in an oxygen-free process.

6) Sintering and aging process

The molded body was sintered at 1010-1150° C. for 4 hours in a vacuum, and then quenched. Next, the obtained sintered body was subjected to two-stage aging treatment at 800° C. for 1 hour and 550° C. for 2.5 hours (both in an Ar protective atmosphere).

<First embodiment>

The alloy shown in Fig. 1 was blended according to the final composition shown in Fig. 2 and Fig. 3, then subjected to hydrogen pulverization treatment, and then finely pulverized by a jet mill to an average particle diameter of 5.0 μm. In addition, the types of alloy raw materials used are also described in FIGS. 2 and 3 . After forming in a magnetic field, it is sintered at 1050°C and 1070°C, and the resulting sintered body is subjected to two-stage aging treatment.

For the obtained R-T-B rare earth permanent magnets, the residual magnetic flux density (Br), coercive force (HcJ) and squareness ratio (Hk/HcJ) were measured with a B-H tracer. Also, Hk is the external magnetic field intensity at which the magnetic flux density becomes 90% of the residual magnetic flux density in the second quadrant of the hysteresis loop. The results are also shown in Fig. 2 and Fig. 3 . 4 is a graph showing the relationship between the amount of Zr added and the magnetic properties when the sintering temperature is 1070°C, and FIG. 5 is a graph showing the relationship between the amount of Zr added and the magnetic properties when the sintering temperature is 1050°C. In addition, the results of measuring the oxygen content in the sintered body are also shown in Fig. 2 and Fig. 3 . In Fig. 2, the oxygen content of Nos. 1 to 14 is in the range of 1000 to 1500 ppm. Also in Fig. 2, the oxygen content of Nos. 15 to 20 is within the range of 1500 to 2000 ppm. In addition, in Fig. 3, the oxygen content of all Nos. 21 to 35 is in the range of 1000 to 1500 ppm.

In Fig. 2, No. 1 is a material not containing Zr. In addition, Nos. 2 to 9 are materials in which Zr is added to low R alloys, and Nos. 10 to 14 are materials in which Zr is added to high R alloys. On the graph of FIG. 4 , a material with Zr added from a low R alloy is represented as "low R alloy added", and a material with Zr added from a high R alloy is represented as "high R alloy added". Also, FIG. 4 is a graph showing the low oxygen content material of 1000 to 1500 ppm in FIG. 2 .

In FIG. 2 and FIG. 4 , the coercive force (HcJ) and the squareness ratio (Hk/HcJ) of the permanent magnet No. 1 sintered at 1070° C. without adding Zr are both low. By observing the structure of the material, it was confirmed that there were coarse grains with abnormal grain growth.

In order to obtain a square ratio (Hk/HcJ) of 95% or more in a permanent magnet added to a high R alloy, it is necessary to add 0.1% by weight of Zr. Abnormal grain growth was confirmed in permanent magnets in which the amount of added Zr was less than this value. Also, as shown in FIG. 6, for example, elemental mapping (mapping) observation by EPMA (Electron Probe MicroAnalyzer: electron probe microanalyzer), B and Zr were observed at the same site, so it is estimated that the ZrB compound was formed. As shown in FIG. 2 and FIG. 4 , when the added amount of Zr increases to 0.2% by weight, the decrease in the residual magnetic flux density (Br) cannot be ignored.

In contrast, the permanent magnet added to a low-R alloy can obtain a squareness ratio (Hk/HcJ) of 95% or more by adding 0.03% by weight of Zr. Also, no abnormal grain growth was found by microstructure observation. Also, even when 0.03% by weight or more of Zr was added, no decrease in remanence (Br) or coercive force (HcJ) was observed. Therefore, according to the permanent magnet added by low R alloy, it is possible to obtain high performance by sintering in a higher temperature range, refining the pulverized particles, and a low oxygen atmosphere. However, even for a permanent magnet added to a low-R alloy, if the amount of Zr added is increased to 0.3% by weight, the residual magnetic flux density (Br) is lower than that of a permanent magnet without Zr added. Therefore, even in the case of a low-R alloy, the addition amount of Zr is preferably 0.25% by weight or less. As with the permanent magnets added to high R alloys, B and Zr cannot be observed at the same site in permanent magnets added to low R alloys, for example, as shown in Figure 7 in EPMA elemental mapping.

When paying attention to the relationship between the oxygen content and the magnetic properties, it can be seen from Fig. 2 and Fig. 3 that when the oxygen content is below 2000ppm, higher magnetic properties are obtained. Furthermore, from the comparison of Nos. 6-8 and Nos. 16-18 in Fig. 2, and the comparison of Nos. 11-12 and Nos. 19-20, it can be seen that when the oxygen content is 1500ppm or less, the coercive force (HcJ) Increase is more preferable.

Next, it can be seen from Fig. 3 and Fig. 5 that the square ratio (Hk/HcJ) of No. 21 without adding Zr is only 86% even when the sintering temperature is 1050°C. It was also confirmed that the permanent magnet had abnormal grain growth in its structure.

For permanent magnets (No.28-30) added with high R alloys, the square ratio (Hk/HcJ) is increased by adding Zr, but the residual magnetic flux density (Br) decreases greatly when the amount of Zr added increases.

In contrast, the permanent magnets (Nos. 22 to 27) added with low-R alloys have improved square ratios (Hk/HcJ) by adding Zr, and have almost no decrease in residual magnetic flux density (Br).

Nos. 31 to 35 in Fig. 3 varied the Al content. It is known from the magnetic properties of these permanent magnets that the coercive force (HcJ) increases by increasing the Al content.

The values of Br+0.1×HcJ are described in FIG. 2 and FIG. 3 . It can be seen that the permanent magnet with Zr added to the low R alloy has a Br+0.1×HcJ value above 15.2 regardless of the amount of Zr added.

For the permanent magnets Nos. 2 to 14 and 16 to 20 in FIG. 2 , the dispersion of Zr in the analysis image was evaluated from the CV value (coefficient of variation) from the EPMA mapping results. Also, the CV value is the quotient (percentage) obtained by dividing the standard deviation of all analysis points by the average value of all analysis points, and the smaller the value, the better the dispersibility. EPMA was JCMA733 manufactured by JEOL Ltd. (PET (pentaerythritol) was used as a spectroscopic crystal), and the measurement conditions were as follows. The results are shown in FIGS. 2 and 8 . From Fig. 2 and Fig. 8, it can be seen that Zr-added permanent magnets (Nos. 2-7) of low-R alloys have better dispersion of Zr than permanent magnets (Nos. 10-14) of high-R alloys added with Zr.

Thus, it can be seen that the good dispersibility obtained by adding Zr to the low-R alloy is the cause of the effect of suppressing abnormal grain growth by adding a small amount of Zr.

Acceleration voltage: 20kV

Irradiation current: 1×10 ^-7 A

Irradiation time: 150msec/point

Measuring point: X→200 points (0.15μm interval)

Y→200 points (0.146μm interval)

Range: 30.0μm×30.0μm

Magnification: 2000 times

<Second embodiment>

Alloy a1, alloy a2, alloy a3 and alloy b1 in Fig. 1 were used to form the final composition shown in Fig. 9, then hydrogen pulverized, and then finely pulverized by a jet mill to an average particle diameter of 4.0 μm. Then it is formed in a magnetic field, sintered at each temperature of 1010-1100°C, and the obtained sintered body is subjected to two-stage aging treatment.

The remanence (Br), coercive force (HcJ) and squareness ratio (Hk/HcJ) of the obtained R-T-B system rare earth permanent magnets were measured with a B-H tracer. Also, the value of Br+0.1×HcJ was obtained. The results are also shown in Fig. 9 . Also, FIG. 10 shows the relationship between the sintering temperature and each magnetic property.

In the second embodiment, in order to obtain high magnetic properties, the oxygen content of the sintered body was reduced to 600-900 ppm by an oxygen-free process, and the average particle diameter of the pulverized powder was made into a fine powder of 4.0 μm. Therefore, abnormal grain growth during sintering is likely to occur. Therefore, the permanent magnets without adding Zr (Nos. 36 to 39 in FIG. 9, shown as Zr-free (Zr-free) in FIG. 10) have extremely low magnetic properties except when sintered at 1030°C. The highest squareness ratio (Hk/HcJ) when sintered at 1030°C was 88%, and did not reach 90%.

Among the magnetic properties, the tendency for the squareness ratio (Hk/HcJ) to be lowered by abnormal grain growth appears first. That is, the squareness ratio (Hk/HcJ) is an index capable of grasping the grain growth tendency. Here, when the sintering temperature range is defined by the sintering temperature range in which the square ratio (Hk/HcJ) of 90% or more is obtained, the sintering temperature range of the permanent magnet without adding Zr is zero.

In contrast, permanent magnets added with low R alloys have comparable sintering temperature ranges. Add Zr 0.05 wt% permanent magnets (Figure 9 No.40~43) and sinter at 1010~1050°C to obtain a square ratio (Hk/HcJ) of more than 90%. That is, the sintering temperature range of the permanent magnet added with 0.05% by weight of Zr was 40°C. Similarly, add Zr 0.08 wt% permanent magnet (Figure 9 No.44～50), add Zr 0.11 wt% permanent magnet (Figure 9 No.51～58) and add Zr 0.15 wt% permanent magnet (Figure 9 Nos.59-66) have a sintering temperature range of 60°C. The sintering temperature range of permanent magnets with Zr 0.18% by weight (Figure 9 No.67~75) is 70°C.

Next, No.37 (sintered at 1030°C, no added Zr), No.39 (sintered at 1060°C, no added Zr), No.43 (sintered at 1060°C, added Zr 0.05% by weight) and No. Fig. 11 shows a micrograph of the cross-section of each permanent magnet of 48 (sintered at 1060° C., adding Zr 0.08% by weight) observed by SEM (scanning electron microscope). Also, the 4πI-H curves of the permanent magnets obtained in the second example are shown in FIG. 12 .

When Zr was not added like No. 37, the crystal grains tended to grow abnormally, and as shown in Fig. 11, some coarse crystal grains were observed. When the sintering temperature rises to 1060°C like No.39, the abnormal grain growth is remarkable. As shown in FIG. 11 , the precipitation of coarse crystal grains of 100 μm or more was conspicuous. No. 43 in which 0.05% by weight of Zr is added can suppress the number of coarse crystal grains as shown in FIG. 11 . Adding Zr 0.08% by weight to No.48, as shown in Figure 11, even after sintering at 1060°C, a fine and uniform structure was obtained, and abnormal grain growth was not observed, and no coarseness of 100 μm or more was observed in the structure. grain.

Next, referring to Fig. 12, compared with the fine and uniform structure of No. 48, when coarse crystal grains of 100 μm or more are generated like No. 43, the square ratio (Hk/HcJ) first decreases. However, no decrease in remanence (Br) or coercive force (HcJ) was observed at this stage. Next, as shown in No. 39, when abnormal grain growth progresses and coarse grains of 100 μm or more increase, the squareness ratio (Hk/HcJ) deteriorates significantly and the coercive force (HcJ) decreases. However, the reduction of the residual magnetic flux density (Br) has not yet started.

The CV values of the permanent magnets NO.51 to 66 in Fig. 9 were measured, and the results are shown in Fig. 9. The CV values are shown in Fig. When it is 100 or less, the dispersibility of Zr is good. However, when the sintering temperature was increased to 1150°C, the CV value exceeded 130 specified in the present invention.

Next, EPMA analysis was performed on the permanent magnet No. 70 in FIG. 9 . FIG. 13 shows scanned images (30 μm×30 μm) of elements such as B, Al, Cu, Zr, Co, Nd, Fe, and Pr. Line analysis was performed on the above-mentioned elements within the range of the scanned image shown in FIG. 13 . Line analysis is performed on 2 different lines. The line analysis distribution diagram of 1 is shown in FIG. 14 , and the analysis distribution diagram of the other 1 line is shown in FIG. 15 .

As shown in FIG. 14 , there are sites (◯) where the peak positions of Zr, Co, and Cu coincide, and sites (Δ, ×) where the peak positions of Zr and Cu coincide. Also, in FIG. 15 , a site (□) where the peak positions of Zr, Co, and Cu coincided was observed. Thus, the Zr-rich region is also a Co and/or Cu-rich region. Also, since the Zr-rich region overlaps with the Nd-rich region and the Fe-poor region, it is known that Zr exists in the grain boundary phase in the permanent magnet.

As described above, in the No. 70 permanent magnet, a grain boundary phase including a region where one or more of Co, Cu, and Nd are enriched with Zr is formed. Also, no sign of a compound formed between Zr and B was observed.

From the analysis of EPMA, the frequency at which the Cu, Co, and Nd-enriched region coincides with the Zr-enriched region was obtained. As a result, it was found that the Cu-enriched region coincided with the Zr-enriched region with a probability of 94%. Similarly, the probability that the Co-rich region coincides with the Zr-rich region is 65.3%, and the probability that the Nd-rich region coincides with the Zr-rich region is 59.2%.

Fig. 16 is a graph showing the relationship between the amount of Zr added, the sintering temperature and the square ratio (Hk/HcJ) in the second embodiment.

It can be seen from FIG. 16 that by adding Zr, in order to widen the sintering temperature and obtain a square ratio (Hk/HcJ) of 90% or more, it is necessary to add 0.03% by weight or more of Zr. It is also known that in order to obtain a squareness ratio (Hk/HcJ) of 95% or more, it is necessary to add 0.08% by weight or more of Zr.

<Third embodiment>

R-T-B series rare earth permanent magnets were obtained in the same process as in the second embodiment except that alloys a1 to a4 and alloy b1 in FIG. 1 were used and blended according to the final composition shown in FIG. 17 . The oxygen content of the permanent magnet was 1000 ppm or less, and no coarse crystal grains of 100 μm or more were observed when the sintered structure was observed. For this permanent magnet, the residual magnetic flux density (Br), the coercive force (HcJ), and the squareness ratio (Hk/HcJ) were measured with a B-H tracer in the same manner as in the first example. Also, the value of Br+0.1×HcJ was obtained, and the results are also shown in FIG. 17 .

The third embodiment was carried out for one of the purposes of confirming the variation of the magnetic properties with the amount of Dy. It can be seen from FIG. 17 that the coercive force (HcJ) increases with the increase in the amount of Dy. On the other hand, all permanent magnets obtained a Br+0.1×HcJ value of 15.4 or more. This shows that the permanent magnet according to the present invention can obtain a high level of residual magnetic flux density (Br) while ensuring a predetermined coercive force (HcJ).

<Fourth embodiment>

Using the alloys a7-a8 and alloys b4-b5 shown in FIG. 1, and blending according to the final composition shown in FIG. 18, R-T-B series rare earth permanent magnets were obtained by the same process as in the second embodiment. In the No. 80 permanent magnet shown in FIG. 18, alloy a7 and alloy b5 were blended at a weight ratio of 80:20. Also, the average particle diameter of the finely pulverized powder was 4.0 μm. The oxygen content of the obtained permanent magnet was 1000 ppm or less as shown in FIG. 18 ; however, no coarse crystal grains of 100 μm or more were observed when observing the structure of the sintered body. For this permanent magnet, the residual magnetic flux density (Br), the coercive force (HcJ), and the squareness ratio (Hk/HcJ) were measured with a B-H tracer in the same manner as in the first example. Also, the value of Br+0.1×HcJ was obtained, and the results are also shown in FIG. 18 .

As shown in Fig. 18, compared with the first to third examples, even when the content of the constituent elements is changed, a higher level of remanence can be obtained while ensuring a predetermined coercive force (HcJ). Flux density (Br).

As described in detail above, by adding Zr, abnormal grain growth during sintering can be suppressed. Therefore, even when processes such as reducing the oxygen content are employed, reduction in the square ratio can be suppressed. In particular, the present invention enables Zr to exist in a sintered body with good dispersibility, so that the amount of Zr for suppressing grain growth can be reduced. Therefore, deterioration of other magnetic properties such as residual magnetic flux density can be suppressed to a minimum. In addition, according to the present invention, since a sintering temperature range of 40° C. or higher can be ensured, even when a large-scale sintering furnace prone to heating temperature unevenness is used, R-T-B-based rare earths having stable and high magnetic properties can be easily obtained. permanent magnet.

Claims (6)

1. A method for manufacturing an RTB system rare earth permanent magnet, wherein the RTB system rare earth permanent magnet is made of a sintered body having the following composition: R: 25 to 35% by weight (R is one or more of the rare earth elements Two or more, but the rare earth element is the concept of containing Y), B: 0.5 to 4.5% by weight, one or two of Al and Cu: 0.02 to 0.6% by weight, Zr: 0.03 to 0.25% by weight, Co: 4 Weight % below but not including 0, and the remaining part is substantially Fe, the method includes: making a Zr-containing low-R alloy containing R ₂ T ₁₄ B compound as the main body, and using R and T (T is Fe or Fe and Co are essential at least one kind of transition metal elements) as a main body of a high-R alloy containing more R than the low-R alloy. The formed body is sintered. 2. The method for producing an R-T-B based rare earth permanent magnet according to claim 1, wherein the low R alloy contains one or two of Cu and Al in addition to Zr. 3. The manufacturing method of the R-T-B series rare earth permanent magnet according to claim 1, wherein the sintering temperature range of the R-T-B series rare earth permanent magnet for obtaining a square ratio Hk/HcJ of more than 90% is above 40°C. 4. The method for producing an R-T-B based rare earth permanent magnet according to claim 1, wherein the Zr content of the sintered body is 0.05 to 0.2% by weight. 5. The method for producing an R-T-B based rare earth permanent magnet according to claim 1, wherein the Zr content of the sintered body is 0.1 to 0.15% by weight. 6. The method for producing an R-T-B based rare earth permanent magnet according to claim 1, wherein the amount of oxygen contained in the sintered body is 2000 ppm or less.

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